Techniques of fabricating all-ceramic dental restorations by hand and methods using commercial high-tech systems such as CAD/CAM systems each have their limitations and target different segments of the dental laboratory market. There are two main challenges restricting widespread use of high-strength ceramic materials for cost-effective fabrication of dental restorations and both challenges are related to the sintering step of the operation. High-strength ceramic materials are crystalline materials formed from powder and require high temperatures for sintering that result in substantial shrinkage. Any technique enabling use of these materials for dental restorations should offer ways to (1) compensate for shrinkage and (2) provide a furnace capable of reaching the temperatures necessary to sinter the material to nearly full density.
A technique reportedly providing the highest strength for manually produced copings, the Vita® In-Ceram™ method (developed by VITA Zahnfabrik), has been advertised as yielding a material with flexural strength of about 500 MPa or even higher. This technique has not become popular mostly due to esthetic limitations and a tedious multi-step fabrication procedure that includes a glass infiltration step. This glass infiltration technique is one way to circumvent the above-mentioned challenges. Vita® In-Ceram™ copings are slip-cast on a gypsum die and soft-sintered with negligible shrinkage. The final glass infiltration step does not require a special furnace. The resulting product is a fully dense restoration having undergone no shrinkage. Nonetheless, the presence of a glass phase in the glass-infiltrated ceramics makes it inferior to corresponding crystalline ceramics in mechanical strength and chemical durability.
Currently available CAD/CAM systems are capable of compensating for shrinkage by milling enlarged shapes. Moreover, high-temperature isotropic sintering results in fully dense and accurate final shapes. However, CAD/CAM systems and procedures are expensive and not affordable by small labs. For example, two of the most recently developed commercial state-of-the art CAD/CAM systems, the LAVA™ system (available from 3M ESPE) and the CERCON® system (available from Dentsply/Degussa), require the purchase of a scanner, milling machine and high-temperature sintering furnace and are currently priced in the range of approximately $60,000-$180,000. Both of the aforementioned CAD/CAM systems employ soft-sintered zirconia blocks. The blocks are milled to enlarged shapes and subsequently sintered to full density. Both systems are advertised as yielding materials having a flexural strength of about 900 MPa or higher.
Glass-ceramic materials obviate the need to compensate for shrinkage and high temperature sintering. They can be hand-built on a refractory die and sintered at fairly low temperatures to assure accuracy of the final shape. One example of such a material is an OPC™ Low Wear (available from Pentron Laboratory Technologies, LLC) porcelain jacket crown (PJC). Glass-ceramic materials can also be injection molded into a refractory investment mold formed by the lost wax technique. Examples of commercially available materials used in this process include OPC® porcelain, and OPC® 3G™ porcelain, IPS Empress® porcelain and Empress 2™ porcelain. The physical mechanism underlying the high processability/formability of these glass-ceramics is the viscous flow of its glass component. The glass-ceramic materials listed above (Optec™, OPC® and OPC® 3G™, Empress® and Empress2™ materials) have from about 40% to about 60% of a glass phase which serves as a matrix in which from about 40% (e.g., Optec) to about 60% of crystals (e.g., Empress2) are embedded. These crystals are grown in-situ by crystallization heat-treatment of the parent glass. Alternatively, in a method described by Hoffman in U.S. Pat. Nos. 5,916,498, 5,849,068 and 6,126,732, in order to improve processability of the material, up to 50% glass is added to the crystalline ceramic powder. As a result, the reported flexure strength is limited to less than 600 MPa. By introducing a glass phase into the microstructure, strength is compromised to gain better processability.
Sintering of glass-ceramic powders is a relatively fast process compared to sintering of crystalline ceramic powders due to the viscous flow mechanism of the former, which is associated with higher densification rates, but the presence of the residual glass phase limits the strength of the final product. Another benefit of the viscous flow mechanism is that the glass ceramic conforms to the shape of the die during sintering without cracking. On the other hand, crystalline ceramics can be much stronger than glass ceramics, but crystalline ceramics sinter by a solid-state diffusion mechanism that is intrinsically slow creating inhomogeneous shrinkage, generating significant sintering stresses that may result in associated cracking. Liquid phase sintering induced by the addition of sintering aids greatly enhances sinterability of crystalline ceramics by promoting particle rearrangement and solution-precipitation mechanisms but such mechanisms do not achieve all the advantages of the viscous flow mechanism.
At the same time many experimental and theoretical studies reveal a decrease of the melting temperature of nanometallic particles in comparison with the melting temperature of conventional bulk metals. Its magnitude depends mostly on particle size and crystal structure as well as particle surface conditions and the host matrix environment such as the presence of impurities, level of agglomeration, coating, deposition substrate and the like. Usually, melting is associated with a pre-melting process resulting in a change in shape of the nanoparticles followed by the formation of a liquid skin on the melting nanoparticles. The liquid skin thickness increases during melting gradually consuming the solid particle core. Transmission electron microscopy studies, such as the one discussed in “Shape Transformation and Surface Melting of Cubic and Tetrahedral Platinum Nanocrystals” by Z. L. Wang, J. M. Petroski, T. C. Green and M. A. El-Sayed, J. Phys. Chem. 102, (32) 6145-6151 (1998), have established that 8 nanometer platinum nanoparticles begin to melt at about 600° to about 650° C., which is a much lower temperature than the melting point of bulk platinum at 1769° C. At about 500° C., cubic particles change their shape to a spherical shape with surface melting occurring at about 600° C. to about 650° C. The molten layer surrounding solid cores of platinum nanocrystals is about 1 nm in thickness at 600° C. and the thickness increases with temperature as the nanoparticles continue to melt. The “melting point depression” abbreviated as MPD is a thermodynamically driven phenomenon and can be explained by a drastic increase in the surface area/volume ratio in nano-particulate materials and the corresponding increase in their specific surface energy. This leads to a size-related dependence of melting temperature that is roughly close to 1/d functionality, where d is the mean particle size, and contains surface tension coefficients, latent heat of melting and the molten skin thickness as parameters.
Table 1 presents some experimental data illustrating the difference in melting temperatures for nanoparticles and the corresponding bulk metals and semiconductors.
TABLE 1Melting temperatures of the selected nanomaterialsNano-MaterialBulkMeltingNano-Nano-MeltingMeltingPointParticleParticleTemperatureTemperatureDepressionMelting/SurfaceMaterialShapesize, nm° C.° C.° C.TMnano/TMbulkMeltingRef.Ptcubic8650176911000.37surface melting[1]Auspherical465010574000.61melting[2]Agspherical74709614900.49melting[2]PdwireDiameter 4.6300144511000.21melting[3]Length 202Snspherical10 153232800.66melting[4]CdSspherical4630167810800.38melting[5]GewireDiameter 556509302800.70surface melting[6]Length 1000from ends[1] “Shape Transformation and Surface Melting of Cubic and Tetrahedral Platinum Nanocrystals,” Z. L. Wang, J. M. Petroski, T. C. Green and M. A. El-Sayed, J. Phys. Chem. 102, (32), 6145-6151 (1998).[2] “Size-Dependent Melting Temperature of Individual Nanometer-Sized Metallic Clusters,” T. Castro, R. Reifenberger, E. Choi and R. P. Andres, Phys. Rev., B 42 (13), 8548-8556 (1990).[3] “Size Controlled Synthesis of Pd Nanowires Using a Mesoporous Silica Template Via Chemical Vapor Infiltration,” K-B Lee, S-M Lee, and J. Cheou, Adv. Mate., 13 (7), 517-520, (2001).[4] “Size-Dependent Melting Properties of Small Tin Particles: Nanocalorimetric Measurement,” S. L. Lai, J. Y. Guo, V. Petrova, G. Ramanath and L. H. Allen, Phys. Rev. Lett., 77(1), 99-102, (1996).[5] “Melting in Semiconductor Nanocrystals,” A. N. Goldstein, C. M. Echer and A. P. Alivisatos, Science, 256, 1425-1427, (1992).[6] “Melting and Welding Semiconductor Nanowires in Nanotubes,” Y. Wu and P. Yang, Adv. Mater., 13 (7), 520-523, (2001).
Onset of surface melting occurs usually at temperatures even lower than the temperature at which the entire nanoparticle melts. It can be speculated that the “molten shells” of the pre-melted nanoparticles work as “a lubricant” inducing higher mobility of the particles and higher diffusion rates and hence facilitating densification at temperatures much lower than 0.6 of the melting point (Tm).
It can be further speculated that thermodynamic considerations explaining the mechanism of MPD described above should hold for ceramic nanoparticles as well. Nevertheless, the MPD effect is not very well studied in ceramics for obvious reasons—even the depressed melting point anticipated for ceramic nanoparticles will still be very high making it extremely difficult to conduct observations similar to those for metals and semiconductors described above in Table 1.
For example, the melting point (TM) for pure alumina and zirconia are 2050° C. and 2700° C., respectively, and therefore the MPD effect of the order of 0.5 TM will result in melting temperatures for nano-alumina and nano-zirconia particles of about 1025° C. and 1350° C. However, there are some indirect indications that MPD does occur in nanoceramics such as extremely low sintering temperatures for nanopowders as reported in R. A. Kimel, Aqueous Synthesis and Processing of Nanosized Yttria Tetragonally Stabilized Zirconia, Ph.D. Thesis, The Pennsylvania State University, the Graduate School, the College of Earth and Mineral Sciences, (2002) and in G. Skandan, H. Hahn, M. Roddy and W. R. Cannon, “Ultrafine-Grained Dense Monoclinic and Tetragonal Zirconia,” J. Am. Ceram. Soc., vol. 77, no. 7, pp. 1706-10 (1994), which are both hereby incorporated by reference.
These studies reported onset of densification at surprisingly low temperatures of about 0.3 TM as well as a surprising and unique ability of nanoceramics to be translucent at fairly high levels of porosity.
Some studies reported extreme difficulty in sintering nanoceramics to full density due to rapid grain growth. For example, Skandan et al. (cited above) observed that grains grew 15 times the initial particle size in the case of nano-zirconia. The other major obstacle encountered with the use of nanoparticles in the fabrication of dental articles is related to difficulties in the consolidating of bulk shapes using conventional methods like powder compaction and slip-casting. It the scope of the present invention to utilize the advantages of nanoparticulate ceramics while successfully overcoming the obstacles currently hampering use of such nanoparticulates as dental ceramics.
It is desirable to provide dental ceramics having low sintering temperatures and high strengths. It would be beneficial to provide dental ceramics having sintering temperatures that are low enough to be sintered in existing dental furnaces, yet maintaining high strength and translucency. It is most desirable to provide processing techniques for dental ceramics that result in fully densified dental ceramics.